Process for the production of high strength steels



June 11, 1968 v. F. ZACKAY ET AL 3,388,011

PROCESS FOR THE PRODUCTION OF HIGH STRENGTH STEELS Filed om. s, 1965 Y3 402,000 psi rs 400,000 psi 40o ,--Ts= 60,000psi "a, 300 YS=303,000psi I .1 i w I f 200 8 Lu I O: I 23 I00- i 1 I l l I o a 2 o l 2 3 4 5 6 7 a 9 10 [2 l3 l4 STRAIN, AUSFORMED H-ll STEEL 40o YS=3I9,000psi TS=322,000psi T5 303,000 psi 72 300 I x i '5 YS= 230,000 psi I I 200 y g m :3 i II I E5 IOO- I I I' o l J L o l 2 3 4 5 6 1 a INVENTORS STRAIN, VICTOR E ZACKAY CONVENTIONAL H-ll STEEL BY EARL R' PARKER ATTORNEY PRGCESE: FGR THE PR ODUCTEQN @F STEEILS Victor F. Zaciray, Berkeley, and Earl R. Parker, Grinds, Qalii, assignors to the United fitates of America as represented by the United fitates Atomic Energy Commission Fiied st. 3, i965, Ser. No. 494,284 ll Qlaim. ((Jl. 148-123) AiidTRACT @F DECLGESURE To produce high strength alloy steel with high ductility, the steel is processed by a series of steps comprised of ternpering, applying stress at an elevated temperature, cooling and post-tempering. The steel is strained at a temperature in the range from 150 r. to 900 F, cooled and then post-tempered. The elevated temperature strain aging results in a substantial increase in the yield strength, and a favorable shape of the stress-strain curve for good mechanical properties.

This invention relates to a method for increasing the yield strength of alloy steel while maintaining a relatively high degree of ductility. The invention described herein was made in the course of, or under, Contract W-740S-ENG-48 with the United States Atomic Energy Commission.

In the normal heat treating of steel, strength is sacrificed to retain high ductility. It has been generally considered that elevated temperature heat treating is incompatible with high strength due to heterogeneous nucleation. Such nucleation tends to occur most easily at lattice imperfections or discontinuities which produce precipitates at grain boundaries and forms sites at which cracks will propagate and spread, causing catastrophic failure of the steel. Therefore, a very fine uniform dispersion of precipitates is needed for very high strength and should come from a homogeneously nucleated precipitation reaction. Thus this invention is directed primarily to those alloy steels having a very fine microstructure. The precipitates in such steels act in several ways as barriers to dislocation movement, such as strong impenetrable non-coherent particles. The precipitates can also act as coherent or incoherent particles through which dislocations can pass, but only at stress levels beyond those required to move dislocations through the parent lattice.

Ductility and toughness in metals depend upon the behavior of dislocations. In a steel suitable for most usages, lastic flow must be able to redistribute a high local load around notches and other discontinuities before cracks can form. In conflict with this requirement of notch ductility is the fact that high strength can only be obtained in ductile materials if plastic flow is inhibited by barriers that restrict the movement of dislocations. Thus, high strength and good toughness are seemingly incompatible requirements and in general the strongest metals have tended to be brittle and ductile ones weak.

The present invention provides for the desirable com bination of properties by a process which includes elevated temperature strain aging of high strength steels which have a high dislocation density and a fine martensitic microstructure. Such strain aging substantially increases States Patent 0 Ice the yield strength, in some cases by a factor as high as 25%. This increase in strength is obtained without the concomitant decrease in ductility which is normal in some tempering operations. There is a slight decrease in duetility, as will be subsequently shown, but this decrease is of a relatively minor nature. In the case of the ausformed steel, ductility is still maintained above 10%. In general, those steels having over 200,000 p.s.i. tensile strength and 10% ductility are usually classified as high strength steels.

It is thought that the strain aging gives a high dislocation density and that the dislocations are pinned by the very fine dispersion. A large number of dislocations is necessary for high strength, and by close control of the temperature during the strain aging a very fine martensitic microstructure is maintained.

Therefore it is an object of this invention to provide steels having a highly desirable combination of properties. 1

A principal object of this invention is to provide a process for increasing the yield strength of an alloy steel while maintaining a substantial degree of ductility.

Another object of the invention is to provide a process whereby an alloy steel can be heat-treated to increase strength and toughness and still maintain ductility.

In general, the invention is applicable to quenched and tempered steels of the class having a fine grained martensitic microstructure. Stress is applied to the steel for a controlled time period, to efifect a plastic strain in the range from about one-fourth percent to about five percent, while the steel is maintained at an elevated temperature but below that at which coalescing of precipitates would occur. The steel is then cooled and may subsequently be tempered if necessary.

In the practice of the invention the exact process conditions will differ for each alloy steel, but can readily be determined experimentally, examples being hereinafter described. Essentially, five parameters are varied in such a Way to yield the optimum results. The -five variables are:

(a) Pre-tempering temperature (b) Amount of strain (c) Strain rate (d) Temperature during strain (e) Post-tempering temperature The pro-tempering and post-tempering are optional steps if the ductility and toughness of the steel is sufficient without tempering. The pro-tempering temperature is generally maintained at about 600 F., the temperature at about which iron carbides form, but in any case below the higher temperature, i.e. about 900 F. at which alloy carbides will form. Thus, the subsequent strain aging at an elevated temperature nucleates the alloy carbides that are still in solution. The amount of strain has an upper limit determined by the plastic limit of the steel and is about 5% for certain steels. The minimum strain which may be employed is the amount sufiicient to form dislocations for nucleation sites to occur and is about for certain other steels. Three differing strain rates 1.35 10- sec., l.35 l0 /sec. and x l0 /sec. have been used and the effect on product strength of the ditfering strain rates was found to be minimal. The temperature of the steel during the strain has a substantial effect on the final properties of the steel and may range from room temperature up to about 900 F. The straining temperature is generally maintained below the post-tempering temperature. By an appropriate selection of the steps and process conditions within the foregoing general limits, the desired combination of strength and ductility is produced.

Two original steels which have been processed in ac difiering strain rates. Comparison of Tables III and IV shows that the strain rate does not appreciably alter the results as compared to the effect of the strain temperatture. However, in each case the yield stress and tensile a stress is substantially increased.

TABLE III [Conventional H-ll steel post-tempered at 900 F., strain rate=1.35 10- /sec.]

Strain Temperature, F 70 300 450 600 750 900 Yield Stress, K p.s.i 330 299 312 328 321 296 Tensile Stress, K p.s.i 330 322 318 328 327 323 Elongation, percent 1. 78 4. 9 4. 4 4. 8 5. 5 7. Reduction in Area, percent 12. 0 10.2 8.8 12. 8 9. 4 14. 0

TABLE IV [Conventional 11-11 steel post-tempered at 900 F., strain rate=1.35 1O- /sec.]

Strain Temperature, F. 70 300 450 600 750 900 Yield Stress, K p.s.i 332 308 326 323 *319 280 Tensile Stress, K p.s 332 316 326 323 322 330 Elongation, percent. 1.9 8. 3 3. 6 2. 7 6. 1 5.8 Reduction in Area, p 15.0 13. 4 12.6 15. 5 14. 0 18. 3

Plotted on Figure 1.

cordance with the invention had the following composition.

The superior properties of the H-ll steel processed by this invention relative to conventional H-ll steel are ap- TABLE I 0 st Mn s P Cr v Mo Conventional H-ll 40 85 .33 015 011 5. 06 5 1. 28 Ausiormed 11-11 .40 1.00 .25 .008 .011 5.00 .5 1.40

The ausformed H-ll steel refers to conventional H-ll 3Q parent from the stress-strain curves on the attached drawsteel which has been processed in accordance with the ing, FIGURE 1. The lower curve is for H-il steel teachings of US. Patent No. 2,934,463 to D. I. Schm-atz quenched and tempered at 900 F. and tested at room et 211., issued Apr. 26, 1960. The ausformed H-ll has a temperature. The upper curve applies to I-I-ll steel, proc finer martensitic niicrostrncture and a high dislocation essed in accordance with this invention, which has been density due to the severe deformation of the ausforin-ing. quenched, temperedv at 900 F, strained 2% at 750 F., The conventional 11-11 was quenched and tempered. post-tempered at 900 F., and subsequently tested at 70 The two steels had the following tensile properties F. As can be seen from the curves the yield stress (YS) prior to the elevated temperature strain aging step of the in reas s from 0 p to 319,000 P- and the present i tio tensile stress (TS) from 303,000 p.s.i. to 322,000 p.s.i. In addition the shape of the curve for the strain aged steel lABLE H is more desirable for engineering applications. Yicld fi B5119 5 Elmlgutitm, Referring to Tables III and IV it is seen that an upper percent limit to the temperature of strain aging conventional Conventional @09 393,090 8 H-11 is indicated by the smaller increase in the yield Ausformed 503mm 360OO0 l4 stress obtained by strain aging at 900 F. When the steel 0 is strain aged at room temperature (70 F.) the yield The steel as processed was in the shape of /2" rolled stress increases, but the ductility is too low to be of strucrods, however, it is obvious that the process is applicable tura'l. use. Thus, it is indicated that the temperature of the to a. wide range of shapes including bar, sheet, red, tube steel during strain aging should be greater than room temand the like. perature.

Referring to Tables III, IV, V, VI and VII, representative results obtained by straining 2% at the specified elevated temperatures are shown. Table III and IV refer to conventional H-ll steel, post-tempered at 900 F. at two Tables V, VI and VIII show representative results obtained by elevated temperature strain aging of ausformed H-11 steels, at three different strain rates and the various indicated tempering temperatures.

TABLE V [Auslorrned II-ll steel post-tempered at 900 IE, strain rate=1.35Xl0* /scc.]

Strain Temperature F 300 450 525 000 750 900 Yield Stress, K 13.5 353 302 359 361 340 311 Tensile stress, K p 381 37 379 377 380 393 Elongation, pereen 11.1 10. 8 12. i 10.8 10.9 8.2

Reduction in Area, p 20. 2 23. 2 1S. 5 10. 0 19. 0 32. 0

TABLE VI [Austorrncd H-11 steel, strain rate=1.35 10- /sec.]

Post-Tampering Temperature, F 900 750 600 500 350 Strain Temperature, F" 300 450 525 (300 300 450 300 300 450 300 450 Yield Stress, K p.s.i 350 362 301 308 399 410 402 391 397 306 Tensile Stress, K p.s.i 378 375 376 398 390 410 402 391 397 306 Elongation, percent... 10.0 9. 5 11.1 5. 35 3. 3 2. 6 0. 5 .7 9. 0 0. 5 Reduction in Area, percent 19. 5 13. 5 26. 0 33. 0 25.0 10.0 21.0 80. 0 30.0 28. 0

TABLE VII [Ausiormed H11.steelpost-tempered at 900 F., strain rate=1.35 10- /sec.]

Strain Temperature, F 300 450 525 600 750 000 Yield Stress, K p.s.i 340 370 390 358 347 Tensile Stress, K ps i 381 374 300 375 386 380 Elongation, percent 8.1 9.6 12.7 11.7 11.0 8.3

Reduction in Area, percent 12. 1 18. 7 27. 0 27. 8 37. 0 40.0

As can be seen from Tables V, VI and VII the yield stress is substantially increased and approaches the values obtained for the tensile stress. The ductility in most cases remains above 10%, which as mentioned previously is considered essential for high strength applications. Although the initial ductility w-as 14%, the decrease in ductility to about 10% is not disadvantageous in view of the large increase in the yield stress.

'In the stress-strain curves obtained from Ausformed H-ll steel, as shown in FIGURE 2, the lower curve shows a yield of 303,000 p.s.i. of Ausform steel tempered at 600 F. The upper curve, for Ausforrn steel processed in accordance with this invention, gives a curve with a better shape and a substantial increase in yield stress; from 303,000 p.s.i. to 402,000 p.s.i. The tensile stress also is increased from 360,000 p.s.i. to 400,000 p.s.i. The increased strength is obtained by strain aging the Ausform steel 2% at 300 F., with subsequent tempering at 600 F.

Considering now a typical example of the steps and process conditions with which the Ausforrn steel data of Table V to VII were obtained, the following procedures were utilized.

Ausform H-ll rods having the properties previously listed in Tables I and II were precision ground to 0113:.001 inch in diameter and 1.1251001 inches in length. The rods were then mounted in a hot testing jig with recording instruments to measure temperature and strain. As noted in Tables V to VIII, three different strain rates were utilized; 1.35 10 /sec., 1.35 10 /sec. and 1.35 X 10 sec. The elongation of the specimen was sensed directly at the end of the specimen by a combination of a mechanical micrometer and a differential transformer. This eliminates possible errors due to thermal expansion of the extension rods.

After straining at the indicated elevated temperature, the specimens were air-cooled by a blower and then given the post-tempering treatment in a furnace controlled within F.

After the post-tempering treatment, the bars were honed to a uniform diameter and finally pulled to destruction at room temperature on an Instron testing :machine, and the tensile data in Tables V to VII obtained.

Three additional steels were tested. Their designation and compositions are shown in Table VIII. These steels were strain aged by packrolling at elevated temperatures. Both Ausform and quenched and tempered conditions of each steel were investigated. The changes in the yield and tensile strengths and the toughness are shown in Table IX.

TABLE V111 C Cr M 0 V Ni Mn Si TABLE IX Strain Temp, Tempering Yield Tensile Elongation, Toughness,

F. Temp, F. Strength, Strength, Percent in in.-lb./in.

K p.s.i. K p.s.i. 1.4

HY-TUF Ausiormed- 302 302 273 277 2. 9 928 302 437 273 276 2. 9 1, 171 302 572 272 272 2. 8 1, 127 302 707 254 255 3. 7 1, 220 437 437 244 261 4. 3 1, 232 437 572 257 263 4. 3 1, 261 437 707 233 243 3. 6 1, 268 572 572 265 267 4. 3 1, 407 572 7 07 246 246 3. 6 1, 449 707 707 242 244 3. 6 1, 511 Unstrained 221 257 5. 0 1, 584 HY-TUF Quencbed 302 302 197 230 6. 4 2, 582 302 437 198 227 5. 7 1, 864 302 572 214 226 5. 7 1, 900 302 707 210 226 5. 7 2, 043 437 437 219 238 5. 7 1, 605 437 572 232 246 6. 4 1, 780 437 707 221 238 5. 0 1, 908 572 572 224 229 4. 3 1, 594 572 707 219 221 5. 0 1, 726 707 7 07 227 231 4. 3 l, 066 Unstrained 186 220 7. 1 1, S81 DGAC Ausiormed 302 302 282 2. 9 379 302 437 276 2. 9 432 302 572 286 2.9 488 302 707 269 270 2. 1 504 437 437 266 266 2. 9 939 437 572 270 270 2. 9 881 437 707 252 853 572 572 248 260 5. 0 1, 104 572 707 234 243 5. 0 1, 190 707 707 234 241 5. 0 1, 047 Unstrained 226 265 5. 7 1, 467 DGAC, Quenched 302 302 192 221 7. 1 1, 978 302 437 191 217 5.0 1, 756 302 572 192 210 6. 4 1, 738 302 707 193 215 6. 4 1, 835 437 437 217 229 4. 3 1, 741 437 572 203 222 5. 0 l, 407 437 707 194 212 5. 7 1, 450 572 572 216 223 3. 6 1, 356 572 707 205 212 4. 3 1, 372 707 707 207 216 5. 0 1, 246 Unstrained 190 217 5. 0 2, 043 51/121, Ausformed 302 302 207 280 5. 0 1, 052 302 437 267 276 3. 6 1, 020 302 572 257 269 5. 0 872 302 707 243 252 5. 7 904 437 437 235 253 5. 0 1, 132 437 572 220 244 5. 0 1, 001 437 707 221 235 6. 4 971 572 572 247 263 4. 3 916 572 707 232 251 5. 0 866 07 707 229 245 5. 7 847 Unstrained 212 256 6. 4 1, 121

TABLE IX--Continued Strain Temp, Tampering Yield Tensile Elongation, Toughness,

F. Temp, F. Strength, Strength, Percent in in.-lb./in.

K p.s.i. K p.s.i. 1.4

5M21, Quenched 302 302 176 214 7. 1 1, 910 302 437 185 211 7. 1 1, 746 302 572 181 206 7. 1 1, 457 302 707 186 209 6. 4 1, 431 437 437 200 209 5. 7 1, 828 437 572 193 206 6. 4 1, 747 437 707 190 202 6. 4 1, 678 572 572 209 219 4. 3 1, 152 572 707 196 208 5. 1, 511 707 707 171 176 7. 1 1, 203 Unstrained 170 212 7. 1 2, 236 Unstrained 172 216 7. 1 1, 966

Attention is particularly invited to the yield and tensile strength of the specimens strain aged at the various strain 1 cracking of the strain aged steel is maintained at a high 2 level.

While the invention has been described with respect to certain specific steels, it is apparent that it is applicable to a wide variety of steels and that many modifications are possible within the spirit and scope of the invention. For instance, it is possible to vary the steel composition to vary the properties obtained. Other methods of straining and differing strain rates at alternate rates may be utilized. Thus it is not intended to limit the invention except as defined by the appended claim.

What is claimed is:

1. A process for increasing the yield strength of a quenched and tempered high alloy steel .mass having a martensitic grain size of about one half the size which would be obtained by the normal quenching of austenite to martensite, comprising:

(a) applying stress to said steel for a controlled length of time to effect a plastic strain of about 4 percent to about 5 percent, while maintaining said steel during said stress within a temperature range of 150 degrees Fahrenheit to about 900 degrees Fahrenheit,

(1)) cooling said steel mass to room temperature, and

(c) post-tempering said steel by heating to a temperature within the range from 350 degrees Fahrenheit to 900 degrees Fahrenheit whereby the yield point is substantially increased.

References Cited UNITED STATES PATENTS 3,230,118 1/1966 Tufts 148l2 2,924,544 2/1960 Nachtrnan.

HYLAND BIZOT, Primary Examiner.

DAVID L. RECK, Examiner.

W. W. STALLARD, Assistant Examiner. 

